Highly Emissive Divalent-Ion-Doped Colloidal CsPb1-xMxBr3 Perovskite Nanocrystals through Cation Exchange

Article abstract of DOI:10.1021/jacs.6b13079

Colloidal CsPbX₃ (X = Br, Cl, and I) perovskite nanocrystals (NCs) have emerged as promising phosphors and solar cell materials due to their remarkable optoelectronic properties. These properties can be tailored by not only controlling the size

Full text of DOI:10.1021/jacs.6b13079

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Article
Highly Emissive Divalent Ion Doped Colloidal CsPb1–
xMxBr3 Perovskite Nanocrystals through Cation Exchange
Ward van der Stam, Jaco J. Geuchies, Thomas Altantzis, Karel H.W. van den Bos, Johannes
D. Meeldijk, Sandra Van Aert, Sara Bals, Daniël Vanmaekelbergh, and Celso de Mello Donegá
J. Am. Chem. Soc., Just Accepted Manuscript • DOI: 10.1021/jacs.6b13079 • Publication Date (Web): 04 Mar 2017
Downloaded from http://pubs.acs.org on March 5, 2017
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Highly Emissive Divalent Ion Doped Colloidal
CsPb1–xMxBr3 Perovskite Nanocrystals through
Cation Exchange
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Ward van der Stam,†║ꢀ Jaco J. Geuchies,†║ Thomas Altantzis,‡ Karel H. W. van den Bos,‡
Johannes D. Meeldijk,§ Sandra Van Aert,‡ Sara Bals,‡ Daniel Vanmaekelbergh,† and Celso de
Mello Donega†*
AUTHOR ADDRESS
†
Condensed Matter and Interfaces, Debye Institute for Nanomaterials Science, Utrecht
University, P.O. Box 80000, 3508 TA Utrecht, The Netherlands
‡
§
EMAT, University of Antwerp, Groenenborgerlaan 171, Bꢁ2020 Antwerp, Belgium
Electron Microscopy Utrecht, Utrecht University, 3584 CH Utrecht, The Netherlands
║ These authors contributed equally to this work
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ABSTRACT
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Colloidal CsPbX3 (X = Br, Cl, I) perovskite nanocrystals (NCs) have emerged as promising
phosphors and solar cell materials due to their remarkable optoelectronic properties. These
properties can be tailored not only by controlling the size and shape of the NCs, but also by postꢁ
synthetic composition tuning through topotactic anion exchange. In contrast, property control by
cation exchange is still underdeveloped for colloidal CsPbX3 NCs. Here, we present a method
that allows partial cation exchange in colloidal CsPbBr3 NCs, whereby Pb2+ is exchanged for
several isovalent cations, resulting in doped CsPb1–xMxBr3 NCs (M= Sn2+, Cd2+, Zn2+, 0<x≤0.1),
with preservation of the original NC shape. The size of the parent NCs is also preserved in the
product NCs, apart from a small (a few %) contraction of the unit cells upon incorporation of the
guest cations. The partial Pb2+ for M2+ exchange leads to a blueꢁshift of the optical spectra, while
maintaining the high photoluminescence quantum yields (>50%), sharp absorption features and
narrow emission of the parent CsPbBr3 NCs. The blueꢁshift in the optical spectra is attributed to
the lattice contraction that accompanies the Pb2+ for M2+ cation exchange, and is observed to
scale linearly with the lattice contraction. This work opens up new possibilities to engineer the
properties of halide perovskite NCs, which are demonstrated to be the only known system up to
date where cation and anion exchange reactions can be sequentially combined while preserving
the original NC shape, resulting in compositionally diverse perovskite NCs.
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KEYWORDS: Perovskite nanocrystals, cation exchange, CsPbBr3, doped nanocrystals
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INTRODUCTION
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The recent development of colloidal CsPbX3 (X = Cl, Br, I) perovskite nanocrystals (NCs) has
resulted in a burst of scientific interest, owing to their outstanding optical properties.1ꢁ3 Due to
their high photoluminescence quantum yields (PLQY’s up to 90%) without any additional
surface passivation, and tunable optical properties throughout the entire visible spectrum,
CsPbX3 NCs are promising new materials for various optoelectronic applications, such as low
threshold oneꢁ and twoꢁphoton pumped4,5 gain materials for lasing and highly efficient LEDs.6
Moreover, recently, colloidal CsPbI3 NCs were used in solution processed photovoltaic cells
with device efficiencies exceeding 10%.7
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Currently, synthetic protocols for colloidal CsPbX3 NCs with a variety of shapes are available,
such as cubes,1,8 nanowires9,10 and nanoplatelets.2,11,12 Besides allꢁinorganic CsPbX3 perovskite
NCs, colloidal organicꢁinorganic perovskite NCs have also been recently prepared.2,13
Furthermore, the composition of colloidal halide perovskite NCs can easily be postꢁsynthetically
tailored through topotactic halideꢁexchange reactions with preservation of the size and shape of
the parent NCs (despite a small lattice expansion or contraction, depending on the size of the
guest halide).14,15 This is in striking contrast to what has been commonly observed for anion
exchange reactions in colloidal IIꢁVI semiconductor NCs, which typically result in severe size
and shape transformations, often leading to hollow NCs.16,17
Topotactic cation exchange reactions are commonplace in colloidal semiconductor IIꢁVI and IIIꢁ
V NCs, allowing access to a variety of compositions and shapes that are not attainable through
direct synthesis protocols.18ꢁ22 In contrast, postꢁsynthetic cation exchange in colloidal CsPbX3
perovskite NCs remains elusive, and has only been addressed in two recent works, which
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reported conflicting observations.14,15 While Akkerman et al. report the observation of CH3NH3
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(MA+) for Cs+ exchange in CsPbBr3 NCs as an undesired complication in anion exchange
protocols using MAꢁX salts,15 Nedelcu et al. state that attempts to exchange either Cs+ or Pb2+
for other cations (Rb+, Ag+, Cu+, Ba2+, Sn2+, Ge2+, or Bi3+) in CsPbBr3 NCs were invariably
unsuccessful and lead only to the decomposition of the parent NCs.14 The resistance of CsPbBr3
NCs to cation exchange was rationalized by Nedelcu et al. as a consequence of the fact that the
perovskite crystal structure is stabilized primarily by the rigid cationic sublattice.14
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The introduction of impurity ions in colloidal IIꢁVI and IIIꢁV semiconductor NCs provides
another synthetic tool to control the optoelectronic properties and to bestow the parent NCs with
novel functionalities, such as magnetism due to unpaired electrons in the dopant,23 or increased
effective Stokes shift due to exciton recombination on the impurity ions.24,25 Furthermore,
impurity doping of perovskite thin films has been shown to improve their performance in solar
cells.26ꢁ28 Recently, Mn2+ doping in colloidal CsPbCl3 perovskite NCs has been achieved by a
direct synthesis method, in which PbCl2 and MnCl2 precursors were mixed in the desired ratio,
leading to NCs with the characteristic Mn2+ emission.29 However, postꢁsynthetic control over the
perovskite composition by introducing different impurity ions has yet to be achieved.
In this work, we present a novel method to perform postꢁsynthetic cation exchange reactions in
CsPbBr3 NCs, through which Pb2+ cations are partially replaced by divalent cations (M = Sn2+,
Cd2+ and Zn2+), yielding doped CsPb1ꢁxMxBr3 NCs. The Pb2+ for M2+ cation exchange results in a
blueꢁshift of the absorption and emission spectra, while preserving the high PLQY’s (>50%) and
narrow PL linewidth (80 meV) of the parent NCs. We attribute the blueꢁshift in the optical
spectra to the contraction of the perovskite cubic unit cell, which results in shorter PbꢁBr bonds
and hence, a stronger ligand field within the Pbꢁhalide octahedra. Interestingly, the blueꢁshift
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scales linearly with the lattice contraction. A blueꢁshift of the PL maximum is also observed
upon anion exchange reactions from iodide to chloride.14,15 However, blueꢁemitting CsPbCl3
NCs obtained by anion exchange methods or direct synthesis protocols have low PLQY’s
(<15%) and broad PL linewidths (>120 meV),1,2,14,15 in contrast to the blueꢁemitting NCs
synthesized with our approach (PLQY >50%, PL linewidth ~80 meV). Our work opens up many
new possibilities to engineer the properties of halide perovskite NCs, which are demonstrated to
be the only known system where topotactic cation and anion exchange can be sequentially
combined, resulting in compositionally diverse colloidal CsPb1–xMxBr3 perovskite NCs (M =
Sn2+, Zn2+, Cd2+ and x ≤ 0.1) with size and shape preservation and high PL quantum yields.
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RESULTS AND DISCUSSION
Optical properties. Colloidal CsPbBr3 NCs (PL maximum at 512 nm) were reacted inside a
nitrogen purged glovebox at room temperature with metal bromide salts (SnBr2, CdBr2, ZnBr2)
and oleylamine dissolved in toluene, which lead to a blueꢁshift of the absorption and PL spectra
(Figure 1). The use of metal bromide salts as precursors for the guest cation ensures that solely
the cation can be exchanged, since both the parent NCs and the precursor contain the same
halide. The spectral position of the PL band blueꢁshifted over a wide range, depending on the
divalent metal cation used and the precursor concentration (Figure 1aꢁc). Interestingly, although
the blueꢁshift of the absorption and PL spectra is observed in all cases (i.e., for all divalent
cations and precursor concentrations), it is more pronounced for specific concentrations and
cations. For example, the blueꢁshift is small (i.e., 30 meV) for both low and high concentrations
of SnBr2 (0.25 mM SnBr2, PL maximum at 506 nm; 1.7 mM SnBr2, PL maximum at 504 nm;
Figure 1d), but is much larger (170 meV) for a SnBr2 concentration of 0.5 mM, resulting in
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efficient PL in the blue region of the visible spectrum (PL maximum at 479 nm, PLQY 62%, see
Supporting Information, Figure S1a). By varying the concentration of SnBr2 in toluene, while
keeping all other reaction parameters constant, the position of the PL maximum can be tuned
between 479 nm and 512 nm. The increase in the SnBr2 concentration also results in the
appearance of a weak and broad absorption feature at ~580 nm. The possible nature of this
transition and of the observed spectral blueꢁshift will be discussed later in this paper.
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Reaction of the CsPbBr3 NCs with other divalent ions (i.e. Cd2+ and Zn2+) leads to more
pronounced blueꢁshifts than reaction with SnBr2, resulting in efficient PL between 452 nm and
512 nm for CsPbBr3 NCs reacted with CdBr2 and between 462 nm and 512 nm for CsPbBr3 NCs
reacted with ZnBr2 (Figure 1e). The key spectral features of the parent CsPbBr3 NCs, such as the
sharp optical absorption edge, wellꢁdefined absorption peaks, and the narrow PL linewidth
(FWHM ~80 meV), are maintained after reaction with metal bromide salts. High PLQY’s over
60% are observed after reaction with divalent cations for all samples (see Supporting
Information, Figure S1a). We note that the PL peak shifts further over time (100ꢁ200 meV over 4
weeks), due to the presence of metal bromide precursor in the stored colloidal dispersions, since
the samples could not be purified due to difficulties with precipitation of the NCs. However, the
NCs do not deteriorate, and maintain their size, shape, high PLQY’s and colloidal stability in
toluene over the course of (at least) several months. Furthermore, we find that the absorption
increases with 2% over the first 90 min of reaction (see Figure S1b, Supporting Information).
The increase in absorption on longer time scales is hard to quantify, since the absorption
spectrum also shifts to higher energies. It has also been observed in other reports that colloidal
perovskite NCs require very high centrifugation speeds to be destabilized from colloidal
suspension and precipitated.30 In the present case, such high centrifugation speeds also lead to
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precipitation of the unreacted metalꢁbromide precursor, while using stronger antiꢁsolvents (e.g.,
methanol) deteriorates the NCs. A recent study by Luther and coꢁworkers reported a novel
purification procedure for perovskite NCs, involving methyl acetate as antiꢁsolvent, which does
not lead to degradation of the NCs, although still requiring high centrifugation speeds.7 We also
used this purification method, but were able to recover only a very small amount of NCs from
the reaction mixtures (see Experimental Section below for details).
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The PL lifetimes of the parent CsPbBr3 NCs and the product perovskite NCs after reaction
with metal bromide salts were measured, displaying multiꢁexponential decay for all samples,
both before and after the cation exchange reactions (Supporting Information, Figure S2).
Therefore, the average lifetime (τavg) was calculated,31 which reveals that τavg is of the same order
of magnitude (~10 ns) in all cases. Furthermore, we note that the higher the PL energy, the
shorter τavg is, as expected based on Fermi’s golden rule.32 For the CsPbBr3 NCs reacted with
SnBr2, the average exciton lifetime (τavg) decreases from 16.8 ns for the parent CsPbBr3 NCs to
6.8 ns for the product NCs obtained upon reaction with 1.7 mM SnBr2. A decrease in τavg is also
observed for the product NCs obtained upon reaction with CdBr2 and ZnBr2, to 6.4 ns and 7.2 ns,
respectively.
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Figure 1. Tunable photoluminescence of CsPbBr3 nanocrystals upon reaction with divalent
cation bromide salts. Photographs of colloidal suspensions under UV illumination of (a)
perovskite NCs after reaction of CsPbBr3 NCs with different concentrations of SnBr2 (0.5 mM,
1.25 mM, 1.7 mM) (b) parent CsPbBr3 NCs (right vial) and product NCs after reaction of
CsPbBr3 NCs with different concentrations of CdBr2 (0.5 mM, 1.0 mM) and (c) parent CsPbBr3
NCs (right vial) and product NCs after reaction of CsPbBr3 NCs with different concentrations of
ZnBr2 (0.5 mM, 1.0 mM). Photoluminescence (full lines) and absorption (dashed lines) spectra
of (d) parent CsPbBr3 NCs (green lines) and product NCs obtained after reaction with different
concentrations of SnBr2 (red and brown lines); and (e) parent CsPbBr3 NCs (green lines) and
product NCs obtained after reaction with different concentrations of CdBr2 (orange lines) and
ZnBr2 (blue lines). In all cases, a blueꢁshift of both the absorption and the photoluminescence
spectra is observed after reacting CsPbBr3 NCs with divalent cation bromides, while the wellꢁ
defined absorption features and the narrow PL FWHM (~80 meV) are preserved.
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Lattice contraction after reaction with MBr2 salts. Transmission Electron Microscopy
(TEM) measurements reveal that the shape and size of the NCs are preserved after reaction with
metalꢁbromide salts (Figure 2 and Supporting Information, Figure S3 and S4), since cubes of ~9
nm are observed in all cases. Electron Diffraction (ED) analysis shows that the atomic reflections
are shifted to larger scattering vectors after the reaction with divalent ions. This indicates that the
atomic CsPbBr3 lattice contracts due to the incorporation of the smaller Sn2+/Cd2+/Zn2+ guest
cations (Figure 2d,h). The 1D Powder ED (PED) patterns (Figure 2d,h) were obtained by
azimuthally integrating the 2D ED patterns, displayed as insets in the TEM images (Figure 2aꢁ
c,eꢁg).33 As displayed in Figure 2, the lattice contraction after reaction with Cd2+ꢁ and Zn2+ꢁ
bromide salts is larger than after reaction with SnBr2. This is expected based on the ionic radii of
the divalent cations, in the case that Pb2+ (r(Pb2+) = 119 pm, coordination number (CN) = 6) is
replaced by the guest divalent cation (r(Cd2+) = 95 pm, r(Zn2+) = 74 pm, r(Sn2+) = 118 pm, CN =
6 in all cases).34 Quantification of the {100} dꢁspacing is displayed in Figure 2, panels d and h.
The diffraction patterns were calibrated with both TlCl and Au reference samples (see
Supporting Information, Supplementary Method 1 and Figure S5 and S6 for details).
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Figure 2. Size and shape preservation and lattice contraction of CsPbBr3 nanocrystals
(NCs) after reaction with MBr2 (reaction time: ~16 h). Transmission Electron Microscopy
(TEM) images of (a) parent CsPbBr3 NCs and (b,c) product NCs obtained after reaction of
CsPbBr3 NCs with different concentrations of SnBr2 precursor (0.5 and 1.7 mM, respectively).
The insets in (aꢁc) display the 2D Electron Diffraction (ED) patterns, which result in (d) 1D
Powder ED (PED) patterns after azimuthal integration (concentration of SnBr2: 0.25, 0.50, 1.0,
1.25, and 1.7 mM, from top to bottom, respectively). The colors of panels aꢁc correspond to the
colors of the 1D PED patterns in panel d. TEM images of (e) parent CsPbBr3 NCs and (f)
product NCs obtained after reaction of CsPbBr3 NCs with ZnBr2 and (g) with CdBr2. The insets
in (eꢁg) display the 2D ED patterns, which result in (h) 1D PED patterns after azimuthal
integration. The colors of panels eꢀg correspond to the colors of the 1D PED patterns in panel h.
A large lattice contraction is observed upon reaction of CsPbBr3 NCs with Cd2+ and Zn2+ (h),
whereas a minor lattice contraction is observed upon reaction with Sn2+ (d). Quantification of the
{100} dꢁspacing is given in panels d and h. Enlarged versions of the TEM images displayed in
panels aꢁc and eꢁg, as well as the TEM images and ED patterns of the other samples, can be
found in the Supporting Information (Figures S3 and S4).
The lattice contraction of CsPbBr3 NCs after reaction with MBr2 salts is also observed in the
highꢁangle annular darkꢁfield scanning transmission electron microscopy (HAADFꢁSTEM)
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images (Figure 3). Here, statistical parameter estimation theory is used to extract structure
parameters from the images.35,36 One of the structure parameters corresponds to the atomic
column positions, which have been used to estimate the absolute distance between neighboring
Pb/halide columns in order to quantify the lattice parameter (Figure 3). By fitting a normal
distribution to these distance estimates, the mean lattice parameter has been determined. The
lattice parameter of the parent CsPbBr3 NCs was calibrated to the value reported in the literature
for the cubic perovskite structure of CsPbBr3. A lattice contraction from |a| = 5.849 ± 0.003 Å to
|a| = 5.839 ± 0.005 Å is observed upon reaction of CsPbBr3 NCs with 0.5 mM SnBr2. The
obtained CsPbBr3 unit vector corresponds well to recent Xꢁray diffraction experiments for the
cubic perovskite crystal structure, but does not exclude the presence of a small portion of
orthorhombic distortions of the octahedra in the atomic lattice.37 Similar analysis on CsPbBr3
NCs reacted with 0.5 mM CdBr2 and 1.0 mM ZnBr2 shows a lattice contraction from |a| = 5.849
± 0.003 Å to |a| = 5.819 ± 0.008 Å (Cd2+) and |a| = 5.808 ± 0.014 Å (Zn2+), respectively. We do
not observe clustered areas with larger/smaller lattice vectors in the images showing the
deviations from the mean lattice vector (Figure 3, middle row, red: larger than mean distances,
blue: smaller than mean distances). From this we deduce that the incorporated guest cations are
homogeneously distributed throughout the NCs. Thickness profiles obtained from the total
scattered intensities of the atomic columns confirm this observation (Figure S7).
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As can be seen, the lattice contraction for Cdꢁ and Znꢁdopants (0.5 % and 0.7 %, respectively)
is larger than for the Snꢁdoped NCs (0.2%), which is in full agreement with the electron
diffraction data presented in Figure 2. The statistical student’s tꢁtest with unequal variances has
been used to verify that the found differences are significant.38 Here, the distribution of the
measured lattice parameters of the parent CsPbBr3 NCs have been compared to a distribution of
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the measurements on the doped CsPb1ꢁxMxBr3 NCs at a significance level of 10%, indicating that
the distributions have different means. Based on these results, we conclude that guest divalent
cations are indeed present in the perovskite NCs and are distributed homogeneously across the
perovskite lattice.
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Figure 3. Quantitative HighꢀAngle Annular DarkꢀField STEM measurements reveal a lattice
contraction of (a) parent CsPbBr3 NCs (|a| = 5.849 ± 0.003 Å) upon reaction with (b) 0.5 mM
SnBr2 (|a| = 5.839 ± 0.005 Å), (c) 0.5 mM CdBr2 (|a| = 5.819 ± 0.008 Å) and (d) 1.0 mM ZnBr2
(|a| = 5.808 ± 0.014 Å) (reaction time: 16 h). In the HAADFꢁSTEM images, the mixed Pb/halide
(bright contrast), Cs (intermediate contrast ) and halide (low contrast) atomic columns can be
easily assigned (top panels). After detecting the positions of the Pb/halide columns from the
HAADFꢁSTEM images, we calculate the interatomic distances (i.e. the lattice spacing, middle
panels). Red bars indicate a larger distance than the mean |a| and blue bars a smaller distance
than the mean |a|. The histogram of distances together with the estimated normal distribution
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reveal a contraction of the cubic perovskite CsPbBr3 atomic lattice upon reaction with SnBr2,
CdBr2 and ZnBr2 NCs (bottom panels).
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Elemental analysis. The results presented above suggest that Pb2+ cations in CsPbBr3 NCs are
replaced by other divalent cations, i.e. Sn2+/Cd2+/Zn2+, resulting in a lattice contraction due to the
smaller ionic radius of the incorporated cations. To verify the incorporation and the distribution
of the guest cations, Energy Dispersive Xꢁray Spectroscopy (EDS) chemical mapping was
performed. We note that the EDS measurements were performed on areas containing a large
amount of NCs, in order to obtain a statistically relevant result. The EDS measurements show
that Cd (Figure 4aꢁe) and Zn (Figure 4fꢁj) are incorporated in the CsPbBr3 NCs in low
concentrations. Moreover, no indication for surface preference or phase segregation is observed,
which leads us to conclude that the distribution of the guest cations within the CsPb1ꢁxMxBr3 NCs
is homogeneous, consistent with the observations presented above. Quantification of the EDS
measurements reveals a Cd:Pb ratio of 0.16 for the product CsPb1ꢁxCdxBr3 NCs obtained after 16
h reaction with 0.5 mM CdBr2 (Supporting Information, Figure S8ꢁS10) and a Zn:Pb ratio of
0.05 for the product CsPb1ꢁxZnxBr3 NCs obtained after 16 h reaction with 1.0 mM ZnBr2
(Supporting Information, Figure S11). The CsPbBr3 NCs reacted with 0.5 mM SnBr2 for 16 h
were also analyzed with bulk EDS measurements, which reveal a Sn:Pb ratio of 0.1 (Supporting
Information, Figure S12). We also performed an EDS quantification of the product CsPb1ꢁ
xCdxBr3 perovskite NCs, and of portions of the TEM grid surrounding clusters of NCs (i.e. the
background). This analysis clearly shows a Cd peak at 3.1 keV for the area containing NCs,
whereas Cdꢁpeaks are below the noise level in the background (see Supporting Information,
Figure S13). The EDS analysis thus clearly demonstrates that divalent guest cations are
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successfully incorporated in CsPbBr3 NCs upon reaction with MBr2 salts. It should be noted that
the elemental concentrations are upper limit estimates, since the samples were not washed prior
to the EDS measurements (see experimental section for details). A few selected samples were
thoroughly purified by precipitating with methyl acetate, which yielded only very small amounts
of NCs (see experimental section for details). The purified NCs showed reduced carbon
contamination upon prolonged exposure to the eꢁbeam (see Supporting Information, Figure S14).
Furthermore, the Sn:Pb ratio for purified NCs (0.085, see Supporting Information Figure S15)
was similar to the Sn:Pb ratio for the unpurified samples (0.1, see Supporting Information Figure
S12). This lead us to conclude that the above presented cation:Pb ratios are valid upper limit
estimates for the incorporated divalent cations. A possible mechanism for the postꢁsynthetic
incorporation of divalent cations into CsPbBr3 NCs will be discussed below.
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It must be noted that we observe bright spots on the CsPbBr3 NCs upon exposure to the
electron beam (see e.g. Figure 4). High resolution HAADFꢁSTEM analysis shows that these
spots correspond to metallic Pb nanoparticles (Supporting Information, Figure S16). This implies
that, upon imaging the NCs with an electron beam, part of the Pb2+ ions are reduced to metallic
Pb. Furthermore, the measured atomic ratios indicate that the parent CsPbBr3 NCs are halide
deficient (Cs:Pb:Br = 1:1:2; ratio cation charge/anion charge= 1.5, Supporting Information,
Figure S8), and that the halide deficiency is preserved after reaction with MBr2 salts (ratio cation
charge/anion charge= 1.5, Supporting Information, Figures S9ꢁS12). Cation/anion ratios
significantly larger than unity are commonly observed for NCs of IIꢁVI (e.g., CdSe) and IVꢁVI
(e.g., PbSe) semiconductors synthesized in the presence of Xꢁtype ligands (e.g., oleate), and have
been rationalized by considering that the stoichiometric NC is coated by a layer of MꢁX units.39
A similar explanation may apply in the present case, since the CsPbBr3 NCs were synthesized
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from Csꢁoleate and Pbꢁoleate. Nevertheless, halide deficiency is commonly observed in bulk
perovskites, due to the very low activation energies for the creation of anion vacancies in these
materials,40ꢁ42 and is likely also present in perovskite NCs. Density Functional Theory
calculations show that anion vacancies do not result in midꢁgap trap states in halide perovskites,
since their energy levels lie within the conduction band.43 Therefore, despite the presence of
halide vacancies, the exciting optoelectronic properties of the parent CsPbBr3 NCs, such as high
PLQY and sharp excitonic features, are preserved in the product CsPb1ꢁxMxBr3 NCs.
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Figure 4. Energy Dispersive Xꢀray Spectroscopy mapping of CsPb1ꢀxCdxBr3 and CsPb1ꢀ
xZnxBr3 nanocrystals. (a) HAADFꢁSTEM image of CsPb1ꢁxCdxBr3 NCs and the corresponding
(b) Cs map, (c) Pb map, (d) Br map and (e) Cd map, demonstrating the presence of Cd in the
perovskite NCs. The inset in panel (a) shows a photograph of a colloidal suspension of the NCs
under UV illumination. (f) HAADFꢁSTEM image of CsPb1ꢁxZnxBr3 NCs and the corresponding
(g) Cs map, (h) Pb map, (i) Br map and (j) Zn map, indicating the presence of Zn in the
perovskite NCs. The inset in panel (f) shows a photograph of a colloidal suspension of the NCs
under UV illumination. The bright spots observed in the HAADFꢁSTEM images are metallic Pb
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nanoparticles, formed upon prolonged exposure to the electron beam (see Supporting
Information, Figure S16).
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Cation exchange mechanism in CsPbBr3 NCs. The results presented above clearly show that
divalent guest cations (Sn2+, Cd2+, Zn2+) can be incorporated into CsPbBr3 NCs by postꢁsynthetic
cation exchange reactions. Considering that aliovalent exchange of Cs+ (radius: 1.88 Å) for M2+
would strongly destabilize the perovskite structure, since it would result in smaller M2+ cations in
sites with coordination number 12 and would require charge compensation, we will assume
below that only isovalent exchange of Pb2+ for M2+ has taken place, leading to partial
substitutional replacement of PbBr6 octahedra by MBr6 octahedra (Figure 5).
Figure 5. Schematic overview of partial cation exchange in CsPbBr3 nanocrystals. Pb2+
cations are partially replaced by other divalent cations (Sn2+, Cd2+, Zn2+) by postꢁsynthetic cation
exchange reactions, resulting in divalent cation doped CsPb1ꢁxMxBr3 NCs (the MBr6 octahedra
are schematically depicted by the blue metalꢁhalide octahedra). Incorporation of smaller divalent
cations results in contraction of the atomic lattice, which induces a blueꢁshift of the optical
transitions (absorption and PL). The halide anions are depicted smaller for clarity.
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Cation exchange reactions have been extensively investigated in NCs of IIꢁVI, IVꢁVI, IIIꢁV,
and IꢁVI semiconductors,18ꢁ22,44,45 and shown to consist of several inherently linked elementary
kinetic steps, which must proceed in a concerted way. The first step is the extraction of the native
cation from the parent NCs, which may occur by a direct place exchange reaction (i.e., extraction
of the native cation and incorporation of the guest cation occur in a direct reaction mediated by a
ligand in solution, e.g. Zn2+ for Cd2+ exchange in ZnSe mediated by oleate ligands44), or by an
independent chemical pathway that is only kinetically coupled to the incorporation of the guest
cation (e.g., cation exchange in Cu2ꢁxS NCs in the presence of excess phosphines).46 The rates of
extraction of the native cation and incorporation of the guest cation must be balanced, otherwise
the accumulation of cation vacancies eventually leads to the collapse of the NC.20,21 There are
several factors involved in the cation extraction process, such as breaking of the cationꢁanion
bonds within the NC and formation of cationꢁligand bonds, while the incorporation of the guest
cation requires breaking the cationꢁligand bonds and formation of the cationꢁanion bonds in the
NC.18,20 The thermodynamic driving force for the cation exchange is determined by the energy
balance of the overall reaction. It should also be noted that the cation exchange process is
essentially a surface reaction, which, in the absence of cation diffusion in the NC, would stop as
soon as all surface cations had been exchanged by the guest cations.44 Therefore, the cation
exchange rates are often limited by the diffusion rates of the outgoing and incoming
cations.18,20,44 Cation diffusion in semiconductor NCs is often attributed to vacancyꢁmediated
migration,18,20,44,45 and its rate is therefore limited by the activation energies for vacancy
formation.
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Cation exchange reactions have not yet been studied in detail in perovskite NCs, but topotactic
anion exchange has been shown to be very efficient and fast in CsPbX3 (X= Cl, Br, I) halide
perovskite NCs,14,15 due to the low activation energy for the formation and diffusion of halide
vacancies in these materials.14,15 In contrast, cation diffusion in perovskites is a very slow
process, owing to the high activation energies for the formation of cation vacancies and the lack
of interstitial sites for interstitial diffusion. A recent study on bulk CH3NH3PbI3 has concluded
that vacancyꢁassisted iodide diffusion in this perovskite material has an activation energy of only
0.58 eV, leading to fast diffusion even at room temperature (diffusion coefficient: 10ꢁ12 cm2.sꢁ1),
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while cation diffusion has much higher activation energies (0.84 eV and 2.31 eV for CH3NH3
and Pb2+, respectively).47 As a result, CH3NH3 diffusion is four orders of magnitude slower than
+
Iodide diffusion, and the Pb2+ sublattice is essentially immobile.47 This implies that cation
exchange processes involving Pb2+ in CsPbX3 perovskites would be limited by very slow
diffusion fluxes for both the outgoing Pb2+ and the incoming M2+ guest cations. Our results show
that this is indeed the case, since the Pb2+ for M2+ exchange reaction is very slow and only occurs
partially, despite the large excess of M2+ cations with respect to the parent CsPbBr3 NCs (from
8×103 to 3×105, see Experimental Methods). The number of Pb2+ cations in a 9 nm CsPbBr3 NC
is ~3600 (unit cell volume: 0.2 nm3, NC volume: 729 nm3). This means that even in the low
precursor concentration regime, there is a twoꢁfold excess of guest M2+ cations with respect to
Pb2+, implying that full exchange would be possible, if a sufficiently strong driving force would
be present. As it will be discussed below, the fact that the cation exchange does not reach
completion despite the large M/Pb ratios used, shows that the driving force is small and that the
reaction may in fact be selfꢁlimited.
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The use of metal bromide precursors has multiple benefits in our cation exchange protocol,
leading to successful partial exchange of Pb2+ for M2+ cations in CsPbBr3 NCs, resulting in
doped CsPb1ꢁxMxBr3 NCs. The presence of Br in the guest cation precursor allows us to take
advantage of the ease of formation of halide vacancies in CsPbX3 NCs and to explore the fact
that CsPbX3 NCs quickly establish an equilibrium with solvated halides in the solution phase,
especially in the presence of amines, allowing fast crossꢁexchange between NCs with different
halides.14,15 Given that MBr2 does not dissociate in toluene, it binds to the surface Br vacancies
as a molecular unit, thereby incorporating MBr2 in the NC. This releases energy due to the
formation of bonds between the incoming Br and the NC, but does not consume energy because
the MꢁBr bonds do not need to be broken. The energy released can then be used to break the
bonds between a PbBr2 unit and the NC. Because the bonds formed and broken are similar, the
energy balance is close to zero, and therefore the surface Pb2+ cations can be quickly exchanged
into guest M2+ cations, despite the low temperatures (i.e. room temperature) used in our cation
exchange protocol. To allow the cation exchange process to continue, two diffusion fluxes have
to be established: inward diffusion of guest M2+ cations and outward diffusion of Pb2+ cations.
As discussed above, these fluxes will be slow, since they have high activation energies, and will
be primarily driven by the increase in entropy that results from the formation of CsPb1ꢁxMxBr3
solid solutions. However, the exchange of Pb2+ by smaller M2+ cations leads to a progressive
contraction of the lattice, which leads to strain within the NC. This increasing strain field
counteracts the entropic gain, eventually becoming sufficiently strong to halt the cation exchange
process. This could be a reason for a selfꢁlimited cation exchange, as we have observed here for
several isovalent ions. A schematic representation of the cation exchange mechanism proposed
above is given in the Supporting Information (Figure S17).
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The fact that the lattice contraction (and therefore the extent of Pb2+ for M2+ exchange) does
not directly scale with the concentration of MBr2 precursors, but instead reaches a maximum at
intermediate concentrations, can be rationalized by considering that the cation exchange is
actually driven by the concentration of halide vacancies in the template CsPbBr3 NCs, as
discussed above. The formation of halide vacancies is promoted by the OLAM molecules present
in solution,14,15 which nevertheless can form OLAMꢁBr complexes with halides from both NCs
and MBr2 precursors. As a result, the increase in the concentration of MBr2 in the reaction
medium will eventually decrease the cation exchange rates by outcompeting the NCs for the
limited supply of OLAM, thereby decreasing the formation rate of halide vacancies.
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We have also performed a series of experiments at fixed SnBr2/NC concentration and total
volume, but different OLAM concentrations (Supporting Information, Figures S18 and S19). We
find that the position of the PL peak after the reaction depends on the amount of OLAM present
in solution, but the exact trend depends on the reaction time: for short reaction times (48 h,
Figure S18) the blueꢁshift is larger for lower OLAM concentrations, while for longer reaction
times (84 h, Figure S19) the opposite trend is observed. These observations are consistent with
the mechanism proposed above and can be rationalized by considering that the cation exchange
is driven by the formation of anion vacancies in the NCs, but is rateꢁlimited by inward diffusion
of the guest cations (i.e., surface sites have to be made available to allow the process to proceed).
Therefore, increasing the concentration of surface halide vacancies does not directly increase the
extent of the cationꢁexchange at shorter reaction times because the cation diffusion rates remain
constant. Under these conditions, OLAM will extract from the NC surface not only Br but also
SnBr2. Therefore, higher OLAM concentrations decrease the relative extent of the cation
exchange in the short reaction time regime by shifting the equilibrium towards solvated SnBr2.
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The inward diffusion of SnBr2 competes with its removal from the surface by OLAM, and, albeit
slowly, gradually moves Sn2+ away from the surface, thereby incorporating it into the NC lattice.
As a result, after sufficiently long reaction times (e.g., 84 h, Figure S19) the concentration of
guest cations incorporated in the NCs will reflect the equilibrium concentration of halide
vacancies, and will thus be larger for higher OLAM concentrations. It should be noted that
higher OLAM concentrations induce shape distortions and structural changes after long reaction
times (Supporting Information, Figure S19), indicating that large excesses of OLAM lead to
partial deterioration of the NCs upon prolonged reaction times, probably due to an imbalance
between the rate of formation of halide vacancies and the cation exchange rates. We have also
observed that the PL intensity of the product CsPb1ꢁxSnxBr3 NCs is higher for reactions carried
out under higher OLAM concentrations (see Supporting Information, Figure S18), which can be
attributed to a better surface passivation, either by OLAM molecules or by oleylammonium
bromide formed inꢁsitu.30
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Lattice contraction and blueꢀshift of the optical transitions. Interestingly, the PL energy of
doped CsPb1ꢁxMxBr3 perovskite NCs scales linearly with the lattice vector (Figure 6)
independently of the type of cation that has been incorporated. This is most evident for the Snꢁ
doped NCs, for which more data points are available (Figure 6a), but is also clear when the other
two cations are taken into account (Figure 6b).
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Figure 6. Photoluminescence (PL) energy and lattice vector correlation in CsPb1ꢀxMxBr3
perovskite NCs. (a) PL energy as a function of the lattice vector in doped CsPb1ꢁxSnxBr3 NCs
obtained by postꢁsynthetic Pb2+ for Sn2+ cation exchange in parent CsPbBr3 NCs. (b) PL energy
as a function of the lattice vector in doped CsPb1ꢁxMxBr3 (M= Sn, Cd, Zn) NCs obtained by postꢁ
synthetic Pb2+ for M2+ cation exchange in parent CsPbBr3 NCs. The different colors of the
symbols in panels (a) and (b) correspond to the colors used to identify the different samples in
Figures 1 and 2.
The bandgaps of ABX3 perovskites are known to increase with the increase of the
electronegativity of both B and X atoms and with the decrease of the unit cell volume.48 It is thus
quite remarkable that the trends for the three cations are very similar despite their different radii,
electronic configurations and properties (Pb2+ [5d106s26p0]: r= 119 pm, electronegativity χ= 1.6;
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Sn2+ [4d105s25p0]: r= 118 pm, χ= 1.7; Cd2+ [4d105s0]: r= 95 pm, χ= 1.5; Zn2+ [3d104s0]: r= 74 pm,
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χ= 1.7; polarizability and Lewis hardness decrease from Pb2+ to Zn2+). This implies that the
observed blueꢁshift of the optical transitions is primarily due to the lattice contraction, and that
the MBr6 guest octahedra are electronically decoupled from the PbBr6 framework. This is
particularly clear for the Snꢁdoped NCs, since it is reported in literature that the band gap of
CsSnBr3 and CsPb1ꢁxSnxBr3 are redꢁshifted with respect to that of CsPbBr3.49,50 This shows that
the bandꢁedge absorption and the PL transition of the CsPb1ꢁxSnxBr3 NCs prepared in this work
are solely determined by the PbBr6 octahedra, despite the presence of SnBr6 groups in the NC.
The additional feature observed at lower energies (540ꢁ600 nm) in the absorption spectra of the
CsPb1ꢁxSnxBr3 NCs (Figure 1d) can thus be ascribed to these electronically isolated SnBr6
octahedra.
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The PL transition of ABX3 halide perovskites has also been reported to blueꢁshift upon
proceeding along the I, Br, Cl series, as clearly illustrated by the PL tunability achieved by
synthetic and postꢁsynthetic compositional control in both bulk and nanocrystalline
ABX3.14,15,51,52 Although this is accompanied by a contraction of the unit cell, the PL shift does
not depend linearly on the lattice contraction, since the widening of the band gap is largely due to
the increase in electronegativity along the halide series, which effectively increases the ionic
character of the BꢁX bond.48,51,52
In fact, our work provides the first clear observation of a linear relationship between lattice
vector and PL energy in perovskites. A recent work reported a blueꢁshift of the PL peak of bulk
CsSnBr3 upon warming (lattice expansion),53 in striking contrast with our results. Similar
observations (i.e., PL blueꢁshift with increasing temperature) have also been reported for
CsPbBr3 NCs54 and bulk CsSnI3.55 However, it should be noted that this behavior (linear increase
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of the bandgap with increasing temperature) is anomalous, since most semiconductors show the
opposite temperatureꢁdependence (i.e., linear increase of the bandgap with decreasing
temperature).55ꢁ58 This anomalous temperatureꢁdependence is also observed for copper halides
(CuBr, CuCl), lead chalcogenides, and black phosphorus, and has been attributed to a dominance
of the lattice thermal expansion contribution over the electronꢁphonon interaction
contribution.55,57 In the specific case of CsSnBr3 and CsPbBr3 the widening of the bandgap upon
lattice expansion is ascribed to simultaneous narrowing and stabilization of both the valence and
conduction bands, as a result of the antibonding character of their maximum and minimum,
respectively.53
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PL shifts have also been recently reported for MAPbX3 bulk films and single crystals subjected
to highꢁpressures.59,60 Interestingly, small redꢁshifts (≤50 meV) were observed in the lowꢁ
pressure regime (below 0.35ꢁ1 GPa), while strongly nonꢁlinear blueꢁshifts (up to 200 meV) were
observed in the highꢁpressure regime (above 1ꢁ2.5 GPa), and were ascribed to amorphization of
the MAPbX3 lattice.59,60 This explanation is however inadequate for the PL blueꢁshifts reported
in our work, since we do not observe any signs of loss of crystallinity, disorder or phase
transitions in the doped CsPb1ꢁxMxBr3 NCs obtained by postꢁsynthetic Pb2+ for M2+ cation
exchange, despite very clear lattice contractions. It should be pointed out that also from this
perspective the behavior of bulk ABX3 halide perovskites seems to be anomalous, since the PL
of most common semiconductors (e.g. GaN, CdTe) blueꢁshifts linearly with increasing
pressure.61,62
The behavior observed for the CsPb1ꢁxMxBr3 NCs prepared in the present work by postꢁ
synthetic Pb2+ for M2+ cation exchange in 9 nm CsPbBr3 NCs is thus in line with the normal
pressureꢁdependence of the bandgap energies of bulk semiconductors, but is opposite to what is
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observed for bulk ABX3 halide perovskites.59,60 This is striking, and is likely due to nanoscale
effects. The contribution of quantum confinement effects to the blueꢁshift observed in the optical
spectra of the CsPbBr3 NCs after the cation exchange reaction is however negligible, since the
NCs under study here are larger than the estimated exciton Bohr diameter for CsPbBr3 (viz., 7
nm),1 and are therefore in the weak confinement regime. Although the reduction in the NC size
should result in an increase in quantum confinement, the effect would be too small to account for
the magnitude of the spectral shifts observed after the cation exchange reactions (up to 300 meV,
which would require a size reduction from 9 nm to ~5 nm, according to the theoretical size
dependence of the band gap reported in ref. 1). Furthermore, Protesescu et al. have shown that
blueꢁemitting CsPbBr3 NCs are in the subꢁ4 nm size regime,1 whereas the CsPb1ꢁxMxBr3 NCs
obtained in this work have a diameter of 9 nm and yet display efficient PL in the blue region of
the electromagnetic spectrum.
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We thus propose that the remarkable behavior observed here for doped CsPb1ꢁxMxBr3 NCs
indicates that they can accommodate larger isotropic compressive strains than bulk ABX3 halide
perovskites, and, as a result, all PbBr6 octahedra within a NC contract without significant
(additional) distortion, so that the compressive strain is homogenously distributed throughout the
NC (as indicated by Figure 3 above). Distortion and tilting of the BX6 octahedra is known to
induce bandgap reduction.48 These distortions are however always accompanied by lattice
constant changes that act in the opposite direction, thereby decreasing the extent of the redꢁ
shift.48 Recent XꢁRay diffraction analysis of CsPbBr3 NCs suggests that orthorhombic distortions
are an important feature of these materials.37 Although the linear relationship observed in the
present work between the degree of lattice contraction and the blueꢁshift of the PL of CsPb1ꢁ
xMxBr3 NCs (Figure 6) does not rule out the possibility of orthorhombic distortions in the NCs, it
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implies that their contribution does not significantly change with the nature of M or the extent of
the replacement, and therefore that the lattice contraction is the dominant effect. The contraction
of the PbBr6 octahedra leads to shorter PbꢁBr bonds and therefore stronger interactions between
the Pb and Br orbitals. This widens the bandgap because the conduction band minimum (CBM)
is composed of the antibonding combinations between Pb(6p) and Br(4p) orbitals,43 and
therefore should shift to higher energies with stronger interactions. Although the valence band
maximum is also antibonding,43 we expect it to shift less than the CBM because it emerges from
the relatively weaker interaction between the Br(4p6) orbitals and the Pb(6s2) orbitals, which are
stabilized by lone pair and relativistic effects.53 In contrast, the Pb(6p)ꢁBr(4p) interaction is
largely responsible for the chemical bonding in APbX3 perovskites as it leads to charge transfer
from the Pb(6p) to the X(np) orbitals.43
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CONCLUSIONS
We present a postꢁsynthetic cation exchange method, which allows us to partially exchange
Pb2+ ions in CsPbBr3 NCs for several other divalent cations (Sn2+, Cd2+, Zn2+). This isovalent
cation exchange results in a blueꢁshift of the PL bands, without loss of the high PLQYs (>60%),
narrow emission bandwidth, and sharp excitonic absorption transitions. The blueꢁshift is
attributed to contraction of the perovskite lattice due to the incorporation of smaller divalent
guest cations (Sn2+/Cd2+/Zn2+), resulting in shorter Pbꢁhalide bonds and hence an increased
interaction between Pb and Br orbitals. We show that the blueꢁshifted PL energy scales linearly
with the lattice vector of the doped CsPb1ꢁxMxBr3 NCs, in striking contrast to what has been
reported for bulk perovskite materials. This is the first example of postꢁsynthetic cation exchange
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in CsPbBr3 perovskite NCs, demonstrating that halide perovskite NCs are the only known system
in which the complete composition can be postꢁsynthetically tailored with size and shape
preservation, by sequentially combining topotactic anion and cation exchanges. This opens up a
library of possible compositions attainable for colloidal CsPbX3 NCs, which might possess
unprecedented and unparalleled optoelectronic properties and may prove beneficial for a number
of applications.
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EXPERIMENTAL METHODS
Materials. Cesium carbonate (Cs2CO3, Sigma Aldrich, 99.9%), 1ꢁoctadecene (ODE, Sigma
Aldrich, 90%), oleic acid (OA, Sigma Aldrich, 90%), Lead bromide (PbBr2, Sigma Aldrich,
99.999%), oleylamine (OLAM, Sigma Aldrich, 70%), Tin(II) bromide (SnBr2, Sigma Aldrich),
Cadmium bromide tetrahydrate (CdBr2.4H2O, Sigma Aldrich, 98%), Zinc bromide (ZnBr2,
Sigma Aldrich, 99.999%), anhydrous toluene (Sigma Aldrich, 99.8%), anhydrous hexane (Sigma
Aldrich, 95%), anhydrous methyl acetate (Sigma Aldrich, 99.5%). OLAM and OA were
degassed at 120 ºC for 1 h prior to use. All others chemicals were used as received.
CsPbBr3 nanocrystal synthesis. The CsPbBr3 NCs were prepared according to the method
described by Protesescu et al.1 First, Csꢁoleate precursor stock solution was prepared. Cs2CO3
(0.814g), 2.5 mL OA and 40 mL ODE were loaded into a 100 mL roundꢁbottom flask. The
mixture was dried under vacuum for approximately 1 h at 120 °C, and then heated under N2 to
150 °C until all Cs2CO3 had reacted with OA. ODE (5 mL) and PbBr2 (0.069g) were loaded into
a separate 25 mL flask and dried under vacuum for 1 h at 120 °C. OLAM (0.5 mL) and OA (0.5
mL) were injected at 120 °C under N2 atmosphere. After PbBr2 had dissolved, the temperature
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was raised to 180 °C and a 0.4 mL portion of the Csꢁoleate stock solution was quickly injected.
We note that the Csꢁoleate stock solution had to be preheated to ~100 °C before injection. After
five to ten seconds the reaction mixture was cooled by an iceꢁwater bath in order to quench the
reaction.
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Purification of the nanocrystals. The CsPbBr3 NCs were purified following the method
described by De Roo et al.30 (for a synthesis based on 69 mg of PbBr2). The crude synthesis
solution was centrifuged for 3 min at 10000 rpm and the colored supernatant was discarded.
Then, 300 ꢂL of hexane was added and the NCs were dispersed using a vortex mixer.
Subsequently, the suspension was again centrifuged for 3 min at 10 000 rpm, after which the
precipitate, containing larger NCs and agglomerates, was discarded. Another 300 ꢂL of hexane
was added to the supernatant, resulting in a colloidal dispersion of CsPbBr3 NCs.
A second purification procedure, based on recent work by Luther and coꢁworkers,7 was also
used for a few selected samples of colloidal perovskite NCs after Pb2+ for Sn2+ cation exchange
(see below). The additional purification of the colloidal CsPb1ꢁxSnxBr3 perovskite NCs was
performed by adding 1 mL of methyl acetate to the crude NC solution, followed by
centrifugation at 13000 rpm for 1 hour. The amount of precipitate was extremely small. After
removing the supernatant, the precipitate was redispersed in 300 ꢀL hexane. Subsequently, 300
ꢀL methyl acetate was added and the NCs were again centrifuged at 13000 rpm. The supernatant
was discarded and the even smaller amount of precipitate was redispersed in 50 ꢀL of hexane
before being dropcast on a Copper TEM grid for further analysis.
Cation exchange. Cation exchange precursor stock solutions were prepared by dissolving 1
mmol metalbromide salt (MBr2, with M = Sn2+, Zn2+, Cd2+) in 10 mL toluene (0.1 M MBr2), in
the presence of 100 µL OLAM. In a typical cation exchange experiment, 1.5 mL diluted NCs in
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toluene (concentration: ~0.01 ꢀM, using an extinction coefficient of 37.9 cmꢁ1 ꢂMꢁ1 at 335 nm,
for cubes of roughly 9 nm)30 and 0.5 mL cationꢁprecursor (different concentrations, ranging
between 0.125 mM and 1.67 mM, [M2+]/[NC] ratio varied between ~8000 and ~300 000) were
mixed and stirred at room temperature for ~16 h. In the majority of the cases, the NCs were not
purified after the cation exchange reactions, due to difficulties with precipitation of the NCs. The
second purification procedure described above, in which methyl acetate is used as antiꢁsolvent,
was applied to a few selected samples after reaction witth SnBr2, but yielded only very small
amounts of NCs, which were only enough to get subꢁmonolayer coverage on a TEM grid.
The influence of the amount of OLAM was investigated by replacing a volume of toluene (5,
10, or 30 µL) by OLAM, while keeping all other variables (i.e., total volume, concentration of
NCs and SnBr2, ratio SnBr2/NCs) unchanged. The products of the reaction were analyzed after
48 hours and 84 hours.
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Transmission Electron Microscopy (TEM) and Energy Dispersive XꢀRay Spectroscopy
(EDS). TEM measurements were performed using a Tecnai20F (FEI) microscope equipped with
a Field Emission Gun and a Gatan 694 CCD camera. The microscope was operated at 200 kV.
EDS measurements were performed on FEI Talos F200X and an aberration corrected ‘cubed’
FEI Titan 60ꢁ300 electron microscope equipped with a ChemiSTEM system63 operated at 200
kV and 300 kV, respectively. Acquisition time for EDS measurements was ~300 s. Samples for
TEM imaging were prepared by dripping a diluted nanocrystal solution in toluene on a carbon
coated polymer film copper grid (300 mesh). The solvent (toluene) was allowed to evaporate
prior to imaging.
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